Molybdenum-base alloy



June 3, 1969 WINSTON H. CHANG 3,447,921

MOLYBDENUM- BASE ALLOY Filed Dec. 21, 1966 mm "G H m 000 mama TEMPE/20 70/95, "F

1TWVTW1TC TI Winsdron H. Chang by 012: His Artt TTWSH United States Patent 3,447,921 MOLYBDENUM-BASE ALLOY Winston H. Chang, Cincinnati, Ohio, assignor to General Electric Company, a corporation of New York Continuation-impart of application Ser. No. 359,188, Apr. 13, 1964. This application Dec. 21, 1966, Ser. No. 609,982

Int. Cl. C22c 27/00; C22f 1/18 US. Cl. 75-176 Claims This application is a continuation-in-part of application Ser. No. 359,188, filed Apr. 13, 1964 now abandoned.

This invention relates to molybdenum-base alloys containing zirconium and carbon, and more particularly to such alloys in which the mechanical properties are controlled by a precipitate of zirconium carbide.

Allowing techniques for the improvement of properties of metals can be divided into two basic types, namely solid solution alloying and dispersed-particle alloying. Although some alloy systems take advantage of both mechanisms, and in some cases one alloying element can serve both purposes, dispersion alloying generally provides the greatest flexibility in terms of amenability to heat treatment for the production of variable properties in the same basic alloy. Of the two modes of dispersion hardening, precipitation hardening is the reversible one in that the other mode, permanent particle dispersion, is not susceptible to heat treatment for the purpose of taking the dispersed phase into solution to soften the alloy for working nor to later precipitation of the dispersed phase in a different morphology.

in the modern design of molybdenum alloys the mechanisms of solid solution alloying and precipitation alloying have been extensively investigated. One of the molybdenum alloying systems on which much work has been done contains a minor but significant amount of titanium, a somewhat smaller amount of zirconium, and some carbon. Several modifications in this alloy system have been investigated and some are presently being commercially exploited. However, in my present investigations concerning the phase relationships and metallurgy of molybdenum alloys which can be precipitation hardened by combinations of titanium, zirconium and molybdenum carbides, I have found that such levels of titanium as have been advocated in the prior art are deleterious to both strength and ductility. Relatively small amounts of zirconium stabilize the precipitated titanium carbides and also form more stable complex carbides of zirconium and titanium. By appropriate small but critical adjustments of the levels of titanium, zirconium and carbon and the ratios of zirconium plus titanium to carbon, improvements over similar prior art molybdenum alloys have been obtained. However, many of the improvements have been in the direction of higher strength at elevated temperature accompanied by lowered workability and ductility at room temperature and below. Previous developments in which zirconium alone is used as a carbide former in the absence of titanium have concentrated on low contents of zirconium or low ratios of zirconium to carbon that have been found in the present investigation to be undesirable.

Furthermore, many of the previously known alloys in the molybdenum-zirconium-carbon system contain significant amounts of Mo C. Interactions between Mo C and other precipitating phases are generally undesirable. Massive deposits of Mo C often are formed as a molten ingot freezes, and the carbon is not readily redistributed through the matrix on decomposition of the Mo C to other precipitates during aging. Also, unstable lower temperature precipitates that convert to Mo C at moderate temperatures do not provide the desired strength levels. Preferential distribution of Mo C. in the grain boundaries is quite deleterious to properties. Thus, molybdenum alloys that can be hardened by precipitation from solution in the absence of diflicult-to-control interactions between different precipitated phases are greatly to be desired, and are not generally found among previously known alloys.

Therefore, it is an object of the present invention to provide molybdenum-base alloys which can be precipitation hardened or solution treated to produce a softer, more workable material, the alloy being capable of developing high strength at elevated temperatures while at the same time retaining substantial and useful ductility at room temperature and below.

Another object of the invention is to provide molybdenum-base alloys with an optimum content of zirconium and carbon such that they can be treated to develop desired properties at elevated temperatures and at room temperature and below.

Another object of the invention is to provide a precipitation hardenable molybdenum-base alloy in which the formation of undesirable M0 0 is avoided.

Still another object of the invention is to provide a precipitation hardenable molybdenum-base alloy which is not deleteriously susceptible to aging during use at elevated temperatures.

Also included as an object of the invention is the provision of zirconium-carbide precipitation-hardenable molybdenum-base alloys which contain a sufiicient amount of zirconium and carbon in preferred proportions to enable a substantial amount of strengthening of the alloy in accompaniment with meeting the above stated objects.

The drawing is a graph illustarting the aging characteristics of any alloy of the invention.

Briefly stated, in one form, the invention provides a molybdenum-base alloy consisting essentially of, by weight, zirconium in an amount of about 1.02.5%, carbon in an amount of about 0.03-0.15 with the atomic ratio of zirconium to carbon being about from 1.5 :1 to 2.5 :1, balance molybdenum, with substantially all of the precipitating phase in said alloy being zirconium carbide. This car-bide may contain some molybdenum in substitutional solid solution, and other elements may be present either in the molybdenum matrix or in the zirconium carbide particles as commercially unavoidable impurities. Some zirconium and some carbon may remain in solid solution in the molybdenum matrix even when the zirconium carbide has extensively precipitated in the alloy, although, for maximum low temperature ductility, it is desirable to have a minimum of carbon left in interstitial solid solution in the molybdenum matrix. Due to embrittlement from dislocation locking by interstitial carbon, the optimum structure for application of alloys of the invention is one in which as much as possible of the carbon content of the alloy is present in the form of precipitated particles of Zirconium carbide which are of as fine a particle size as can be obtained and are randomly distributed throughout the molybdenum matrix without directional orientation or preferential distribution in the grain boundaries. Other examples of alloys of the invention are those containing 1.0-2.0% zirconium, 0.03-0.1% carbon and have a zirconium-to-carbon atomic ratio between 1.5:1 and 2.5 :1; 1.02.0% zirconium, 0.030.1% carbon, having a zirconium-to-carbon atomic ratio between 2.0: and 2.5: 1; 1.l- 1.3% zirconium, 0.060.07 carbon, having a zirconium-tocarbon atomic ratio between 2.3: l and 2.5 l; and the alloy containing about 1.18% zirconium, about 0.067% carbon, and having a zirconium-to-carbon atomic ratio of about 2.4:1. Percentages in this specification are given in weight except where indicated otherwise.

Even though the zirconium-carbide dispersoid is much harder than the molybdenum matrix, it also is desirable that the zirconium carbide particles be of a size small enough that there be few if any dislocations in each particle, and so that dislocation movement in the particle is very limited if existent at all, thereby greatly increasing the effective hardness of the dispersed particles. Such small particle sizes also benefit the properties of the alloys by decreasing the interparticle spacing.

Relative to zinconium-carbide precipitates in alloys of the invention, molybdenum carbide generally precipitates in the form of massive, sometimes feathery particles due to the excessive Mo/C ratio. They serve little useful purpose in strengthening the alloy or in making it ductile. Therefore, it is quite desirable that the carbon present in the alloy be utilized in the formation of zirconium carbide rather than being trapped in the useless or deleterious molybdenum-carbide form. Although molybdenum carbide serves a purpose in some related alloys as an intermediate phase for the purpose of promoting controlled aging reactions, it is not generally desirable as a constituent of a final alloy in condition for application. Moreover, residual molybdenum carbide left after controlled heat treatment in processing can cause deleterious onthe-job aging and embrittling of such alloys by conversion to other stronger carbides at elevated temperature during application, particularly with the aid of straininduced aging. Although for some applications it may be desirable to have an alloy that will age during use, it is also quite desirable to have alloys that will not age during use for other applications.

I have discovered that, if the zirconium and carbon contents and their relative ratios are kept within certain critical limits, substantially all of the precipitate formed will be zirconium carbide, with the essential or complete elimination of molybdenum carbide. Therefore, the alloy of my invention satisfies the needs and avoids the problem expressed in the preceding paragraph. Although some of my experiments were designed to study the relative reactions of precipitation of molybdenum carbide and zirconium carbide, the alloys of the present invention were discovered to have compositions which do not allow the formation of molybdenum carbide or its interaction with zirconium carbide. The precipitation phenomena are governed by the nature and stability ranges of the carbides, aging kinetics, carbide morphology and other thermodynamic relations which interact in their complexities to produce alloys having specific combinations of strength and ductility which can be varied through thermal and mechanical treatments. I have discovered that the zirconium and carbon contents and their relative proportions in alloys of the present invention combine in their effects on the above and other criteria to produce highly desirable properties and combinations of properties.

As to the limits on composition and proportion, alloys with excessive ratios of zirconium to carbon are either too low in carbon or too high in zirconium to have the desired properties due to deficiency in carbide dispersion in the former case and uncontrollable aging reactions, excessive carbide size, formation of intermetallic phases, or excessive depression of melting temperature in the latter. Similarly, alloys with insufficient zirconium to carbon ratios abound in massive moylbdenum carbide, whose presently unwanted characteristics have been dscussed above. Moreover, too low a content of zirconium has a doubly undesirable efiect of (1) not providing a sufficient quantity of zirconium carbide precipitate for the desired purposes and (2) leaving too much carbon in interstitial solid solution in the molybdenum metal matrix, thus worsening low temperature brittleness of the alloy. For these reasons, the alloy composition and proportion limits stated above and in the claims are critical to the usefulness and desirability of alloys of the invention.

As an example of the alloy outside the scope of the present invention, an alloy composition of 1.66% zirconium .04% carbon, less than 0.005% oxygen, balance molybdenum, has an atomic ratio of zirconium to carbon of about 5.47 (to 1). Such a ratio is excessive for my purpose and would result in uncontrollable aging reactions, excessive solid solution hardening, depression of melting temperature or combinations of these factors which may be acceptable for certain purposes but would be harmful to the uses for which my alloys have been designed. Excessive solid solution hardening is a major cause of embrittlement in some molybdenum alloy systems. A second example outside the invention is an alloy having a zirconium content of 0.72% and a carbon content of 0.03%. Such an alloy would have an atomic ratio of zirconium to carbon of 3.08. Although that ration is within the upper limit of the range of ratios allowed in the alloys of my invention, the zirconium content is too low for the alloy to be used in the manner in which alloys of the present invention can be used, as discussed above.

Another alloy of the prior art has a composition of 0.49% zirconium, 0.022% carbon, balance molybdenum, with a zirconium-to-carbon atomic ratio of 3.311. Tensile test data for this alloy, referred to as Alloy S, are compared in Table I below with data for an alloy of the invention which has a composition of 1.18% zirconium, 0.065% carbon, balance molybdenum, having a zirconium-to-carbon atomic ratio of 2.411, and which is referred to as Mo-S.

TABLE I.COMPARISON OF ALLOYS S AND Mo-5 Test K s.i. Percent Temp. Test Alloy F.) Condition U'lS YS EI RA S SR 118. 7 80. 8 33 56 RX 85. 0 (i3. 4 49 50 Mo-5 75 SR 137. 8 121. 0 24 53 RX 82. O 62. 0 0 5 S 1, 800 SR 72. 7 62. 8 16 34 RX 47. 0 10 88 Mo-5 2,000 SH 82. 2 77. 1 10 67 RX 51. 0 34. 6 29 82 S 2, 400 SR 26. 6 24. 0 27 25 RX 23. 0 10. 9 81 93 Mo5 2, 500 SR 46. 0 40. 3 46 86 RX 33. 3 25. 6 4'.) 86

It can be seen from the table that there is a substantial increase in both ultimate tensile strength and yield strength of Mo-5 over Alloy S. At room temperature, this is an increase in ultimate tensile strength of about 16% and yield strength of almost 51%. This dilference in strength is due in substantial part to the difference in Zr/C ratio, although it is also due in part to other differences in composition. It should be realized that all the examples of prior art differ from Alloy Mo-S both in composition and in ratio, and ratio cannot be varied without also changing composition.

At elevated temperatures, Alloy S is compared at 1800 F. with Mo-S at 2000 F. Mo-5 is shown to about 14% stronger at the 200 F. higher temperature. The difference in strength would be much greater in favor of M-S if both alloys were tested at the same temperature of 1800 F. Comparing Alloy S at 2400 F. with Alloy Mo-S at 25 00 F., it is seen that Alloy Mo-S is about 73% stronger at the F. higher temperature. The same relationship applies, and Alloy Mo-S would be stronger than Alloy S by an even larger percentage when measured at the same temperature.

EXAMPLES As examples of a presently preferred method for producing an alloy of the invention, the following detailed description is given along with information on the properties and characteristics of the alloy so produced.

An electrode weighing about 21 pounds and being about 1.98 inches in diameter and 24.5 inches in length was prepared by hydrostatically pressing two equal lengths of a blended powder containing 1.3% zirconium, 0.15% carbon, balance molybdenum, vacuum sintering at 3000 F. for one hour, and then mechanically joining the two pieces together. The materials used to produce the powder lot pressed into the electrode were as follows: the molybdenum was of a -325 mesh particle size, containing 0.0920% oxygen and 0.002% carbon; the zirconium had a particle size of 200 mesh, and contained 0.145% oxygen, 0.009% nitrogen and 1.8 hydrogen; and the carbon was Grade 38 powder procured from the Acheson Colloids Company of Port Huron, Mich. The electrode was consumably arc melted in vacuum at a maximum pressure of 60 microns and a ratio of mold to electrode diameter of 1.8 to form an ingot 3.5 inches in diameter and 5 inches in length of quite satisfactory quality. Direct current was used with a yield of over 90% of the electrode weight being incorporated into the ingot. The analyzed composition of the ingot was 1.18% zirconium, 0.065% carbon,'0.0018% oxygen, 0.0016% nitrogen and 0.006% hydrogen, corresponding to an atomic ratio of zirconium-to-carbon of 2.411.

After cutting oh the bottom one-inch section to assure complete removal of the starting material, the ingot was machined into a billet 3 inches in diameter by 4.75 inches in length with a beveled nose in preparation for extrusion. During the machining operation, the soundness of the ingot was confirmed as no porosity was observed. The machining chips of the last two cuts in the lathe turning operation were used for the chemical analysis, the results of which are presented above.

In studies of the annealing characteristics of the cast alloy, it was found that temperatures of about 3500* F. were necessary to cause extensive carbide dissolution during the one-hour anneal. Such dissolution became quite evident after annealing at 3750 F. Based on the heattreatment results, a primary processing temperature of 3500 F.-3650 F. was deemed desirable to achieve carbide dissolution and processing ease without excessive grain growth.

For extrusion, the billet was held at 3650 F. for minutes and then extruded to a one-inch diameter, correponding to an extrusion ratio of 9 to 1. A force of 6 10 tons was required for the extrusion. The extruded bar gave a fully satisfactory appearance and produced useful material equivalent to about an 87% yield from the ingot. The extruded condition had a fine-grained structure which was about 6 5% recrystallization with hardnesses of 221 and 229 Vickers (VHN) in the longitudinal and traverse directions, respectively.

Approximately 4 inches of the extruded material were swaged to 0.25 diameter bar stock using a procedure which comprised swaging at 2500 F. for approximately 60% reduction of area, and lowering the temperature to 2100" -F. for final swaging to a total of 93% reduction of area.

Hardness changes in the swaged material upon onehour annealing at various temperatures between 1500" F. and 4000 F. are shown in the drawing and in combination with the approximate grain sizes in millimeters for the recrystallized specimens in Table II below. Rapid softening occurred at 2500 F. coincidental with incipient recrystallization. Complete recrystallization, however, did not occur until 3250" F. at which temperature spheroidization of the zirconium-carbide phase became plainly evident. The alloy showed high resistance to grain growth up to 3750 F. A five-fold increase in grain size, however, occurred between 3750 F. and 4000 F. Using small specimens and moderately rapid cooling by filling the vacuum furnace with helium gas, there was a tendency toward hardening upon annealing at higher temperatures.

TABLE II.EFFECT OF ONE-HOUR ANNEALING ON HARDNESS OF SWAGED ALLOY Annealing Hardness, Grain Size, Temp, F. VHN mm.

As swaged TABLE TIL-EFFECT OF AGING ON HARD- NESS 0F SWAGED ALLOY Aging Annealing Hardness, Temp., F. Temp., F. Time, hrs. VHN

The dispersed phases in representative conditions of the alloy (as-cast, as swaged, annealed at 3750 F., and aged at 2500.F.) were extracted and subjected to X-ray diffraction analysis. The diffraction data and lattice parameter values indicates that, with the possible exception of a very small and insignificant amount of ZrO the only dispersed phase was the face-centered cubic ZrC with a lattice parameter of a equal to 4.657-4.687 angstrorn units. This is somewhat less than the lattice parameter of ZrC (4.699 angstroms) indicating some substitution of molybdenum for zirconium. The elimination of the Mo C phase excluded age hardening by virtue of the Mo C ZrC reaction. This contributed to the lack of pronounced aging response in the alloy. The atomic ratio of zirconium to carbon of 2.4:1 exceeded the threshold ratio below which Mo C could co-exist with ZrC.

Tensile properties of the alloy were obtained in several stress-relieved conditions and in annealed and subsequently aged conditions at several temperatures between room temperature and 3500 F. In addition, tensile tests Were conducted on specimens stress-relieved at 2100 F. for one hour at temperatures between room temperature and F. The results are presented in Table IV below in which K s.i. stands for thousands of pounds per square inch, SR stands for stress relieved and RX means recrystallized. All of the tests were conducted on an Instron machine, using a button-head type specimen with a 0.16 inch diameter and a one inch gauge section. The nominal strain rate was 0.05 per minute. The yield strength was obtained by intersecting the loadtime curve at 0.2% offset. Tests above room temperature were conducted in a vacuum of about 10- torr.

TABLE IV.TENSILE PROPERTIES OF SWAG ED ALLOY K s.i. Percent Temp., One-Hour Heat Tensile Yield Elonga- Reduc- F. Treatment Strength. tion tion in Area -120..- (A) SR 2,100 F 157.0 147.0 5. 5 5.9

100 (A) SR 2,100 F 160.0 151. 11. 8 '10.

(A) SR 2, 100 F.-- 157.0 151.0 9.0 9. 5

72.... (A) SR 2,100 F.-- 156.5 144.5 24.1 47.5

25-... (A) SR 2,,100 F--- 157.8 135.2 18.3 '16. 9

0 (A) SR 2,100 F. 143. 3 128. 3 16. 7 14. 4

SR 2, 100 F- 141. 4 121. 5 18. 0 15. 0

SR 2, 100 F 137. 8 121.0 23.5 52. 7

) SR 2, 500 F. 121. 0 87.0 26.7 51. 8

(C) SR 2, 750 F 112.0 74. 6 13.0 11. 9

(D) SR 3, 000 F. 81.0 44.3 7.3 5.0

(E) RX 3, 250 F 82. 0 62. 0 6. 3 "4. 5

(F) Annealed 3, 750 F 60.2 51.0 1. 5 "1. 0

(H) F+2,750 F. 71.9 33. 6 30. 6 59. 7

(I) F+3,000 F 60. 8 31. 2 30. 8 60. 9

1,000.-.. (A) SR 2,100 F 100. 2 85.9 14. 3 67. 3

( SR 2,500 F-.- 82. 5 61. 4 21. 8 72. 7

0) SR 2,750 F.-- 77.9 51.3 27.8 72.6

(D) SR 3,000 F- 61.0 24. 4 32. 9 79. 1

(E) RX 3,250 F 64. 0 39. 0 33. 5 76.5

(F) Annealed 3,750 F. 61. 9 27. 7 29. 8 73. 8

(G) F+2,500 F 63. 7 27. 9 31. 5 75. 7

(H) F+2,750 F 63. 5 31. 1 32. 0 73. 1

(I) F+3,000 F 61. 1 27. 2 33.0 75. 5

1,500.-.. (A) SR 2,100 F. 00. 7 82. 2 14. 0 70.8

(B) SR 2500 F. 76. 7 59. 9 20. 0 71. 8

(C) SR 2,750 F. 6!). 5 46.0 21.6 76.1

(1)) SR 3,000 F 54. l) 22. 4 25. 0 78. 7

(E) RX 3,250" F 56. 3 38. 1 28. 0 81. 7

(F) Annealed 3,750 F. 66. 3 35. 2 22. 5 75. 8

(F) Annealed 3,7 79. 2 35. 9 17. 9 69. 3

(G) F+2,500 F.. 75. 1 32. 4 18. 1 70.7

(I) F+3,000 F 49. 7 27. 2 32. 3 80. 4

2,500.... (A) SR 2,100 F 46. 0 40. 3 46.0 86. 0

(13) SR 2,500 F.. 42. 2 38. 0 37. 2 84. 7

(0) SR 2,750 F.. 36. 3 32.7 42. 3 87. 0

(D) SR 3,000 F. 29. 7 18. 5 65. 0 86. 2

3,000--.. (A) SR 2,100 F 20. 3 19. 1 46. 1 82. 7

(B) SR 2,500 F 19. 1 17. 3 51.2 82. 1

Radius failures.

The stress-relieved condition is seen to have an excellent low temperature ductility, with reduction in area at 100 F. and 50% reduction in area at room temperature. The high temperature strength of the alloy in this condition was about 20,000 p.s.i. This decreased to about 10,000 p.s.i. at 3500 F. which exceeded the recrystallization temperature by some 250 F. Corresponding to incipient recrystallization, a noticeable reduction in strength of about 50% to 46,000 p.s.i. was

also noted when the test temperature was raised from 2000 F. to 2500 F.

In contrast to the stress-relieved condition, the recrystallized condition was less ductile with about 5% reduction in area at room temperature accompanied by failure in the radius of the specimen. The serrated characteristics of the stress-strain curves at 1500 F. and 2000 F. suggested that, upon recrystallization at 3250 F., sufficient carbide was dissolved to impair the ductility of the alloy at room temperature. However, the amount of carbon supersaturation was apparently insufficient to result in pronounced strain-induced precipitation, as evidenced by the lack of an inverse tempenature dependency of strength at the intermediate temperatures. Decreased ductility in the recrystallized condition is a common phenomenon in metals, particularly in refractory metals, and may be due in part to grain boundary segregation, substructures in the grains, and increased interstitially dissolved carbon. Some improvements in ductility and strength of the recrystallized condition can be seen in the results of the various postrecrystallization anneals presented in Table IV.

It is contemplated by the appended claims to cover any such modifications as fall within the true spirit and scope of this invention.

What I claim as new and desire to secure by Letters Patent of the United States is:

1. A molybdenum-base alloy having a microstructure consisting essentially of a molybdenum matrix having precipitated carbides distributed therein wherein said precipitated carbides are substantially all zirconium carbide, said alloy consisting essentially of, by weight, about 1.02.5% zirconium, about 0.03-0.15% carbon, balance molybdenum, with the atomic ratio of zirconium-to-carbon being about between 1.5 :1 and 2.5 :1.

2. A molybdenum-base alloy according to claim 1 consisting essentially of, by weight, about 1.0-2.0% zirconium, about 0.03-0.1% carbon, balance molybdenum, with the atomic ratio of zirconium-to-carbon about between 1.521 and 2.521.

3. A molybdenum-base alloy according to claim 1 consisting essentially of, by weight, about 1.0-2.0% zirconium, about 0.03-0.1% carbon, balance molybdenum with the atomic ratio of zirconium-to-carbon being about between 2.021 and 2.521.

4. A molybdenum-base alloy according to claim 1 consisting essentially of, by weight, about 1.11.3% zirconium, about 0.060.07% carbon, balance molybdenum, with the atomic ratio of zirconium-to-carbon being about between 2.311 and 2.5:1.

5. A molybdenum-base alloy according to claim 1 consisting essentially of, by weight, about 1.18% zirconium, of about 0.067% carbon, balance molybdenum, with the atomic ratio of zirconium-to-carbon being about 2.4:1.

References Cited UNITED STATES PATENTS 2,678,271 5/ 1954 Ham et al. -176 2,960,403 11/ 1960 Timmons et al. 75-176 3,169,860 2/ 1965 Semchyshen 75-176 3,194,697 7/1965 Chang 75-176 OTHER REFERENCES Nb, Ta, Mo & W, Elsevier Publishing Co., New York, 1961, relied on pages 322 and 323.

WADC TR 59-280, Development of Molybdenum- Base Alloys, Semchyshen et al., October 1959, pages 51-55.

CHARLES N. LOVELL, Primary Examiner.

US. Cl. X.R. 

1. A MOLYBDENUM-BASE ALLOY HAVING A MICROSTRUCTURE CONSISTING ESSENTIALLY OF A MOLYBDENUM MATRIX HAVING PRECIPITATED CARBIDES DISTRIBUTED THEREIN WHEREIN SAID PRECIPITATED CARBIDES ARE SUBSTANTIALLY ALL ZIRCONIUM CARBIDE, SAID ALLOY CONSISTING ESSENTIALLY OF, BY WEIGHT, ABOUT 1.0-2.5% ZIRCONIUM, ABOUT 0.03-0.15% CARBON, BALANCE MOLYBDENUM, WITH THE ATOMIC RATIO OF ZIRCONIUM-TO-CARBON BEING ABOUT BETWEEN 1.5:1 AND 2.5:1. 